Alloying Pre-alloyed Fe–Mo Powders by Silicon Carbide Addition

Nattaya Tosangthum; Thanyaporn Yotkaew; Rungtip Krataitong; Monnapas Morakotjinda; Preeya Nakornkaew; Piyanuch Nakpong; Ruangdaj Tongsri

doi:10.2355/isijinternational.ISIJINT-2024-139

Abstract

This work demonstrated that the silicon carbide addition to pre-alloyed Fe–Mo powder could result either in the formation of steel or iron microstructure depending on the added silicon carbide content. With 1.0 wt.% silicon carbide addition, slowly cooled sintered Fe–Mo–Si–C alloys showed steel microstructures consisting of proeutectoid ferrite and eutectoid transformation product in the form of ferrite + carbide mixture. With 2.0 wt.% silicon carbide addition, slowly cooled sintered Fe–Mo–Si–C alloys with Mo contents of ≥ 0.85 wt.% microstructures comprised ferrite + austenite constituents in the forms of either degenerate upper bainite or ausferrite. With ≥ 3.0 wt.% silicon carbide addition, ductile iron-like microstructures were developed in sintered Fe–Mo–Si–C alloys. The change of microstructures in experimental sintered alloys was attributed to the combined effect of alloying molybdenum, silicon, and carbon elements. Tensile strength and hardness increased with increasing added SiC content while ductility varied with microstructural components.

1. Introduction

The role of alloying silicon (Si) element in steels is one of the most interesting issues of steel metallurgy. The well-known function of Si is ferrite stabilization. When it is dissolved in ferrite, solid solution strengthening by Si can be achieved.¹⁾ The Si in solid solution in ferrite in pearlitic PM steels also results in strengthening.²⁾ Apart from solid solution strengthening, Si also acts as carbide precipitation inhibitor. With proper amount, Si can hinder carbide precipitation in austenite.³⁾ The carbide precipitation retardation is attributed to extremely low solubility of Si in cementite. The dissolution of Si in cementite leads to thermodynamically unfavored formation energy of cementite.⁴⁾ Thus, judicious Si amounts are added to produce advanced carbide-free bainitic (CFB) steels.⁵⁾ During phase transformation step of CFB steels, the carbon (C) partitioning from bainitic ferrite plates/laths to adjacent untransformed austenite plates/laths results in C enrichment of austenite. With carbide precipitation retardation by Si the microstructures of CFB steels comprise bainitic ferrite and C-enriched austenite plates/laths. With high C concentration, carbide precipitation retardation by Si is not strong due to large driving force for carbide precipitation.⁶⁾ Lamellar pearlite can be formed in Si-containing high C steels⁷⁾ and cast irons.⁸⁾

Lamellar pearlite can also be formed in sintered high Si high C alloys. Sintered Fe–Mo–Si–C alloys, produced from mixtures of pre-alloyed Fe–Mo powders and fixed 4.0 wt.% silicon carbide (SiC), showed microstructures resembling that of a ferritic-pearlitic ductile iron (DI), i.e., the microstructural feature comprised a black particle enveloped with ferrite halo, pearlite and ausferrite.⁹⁾ The lamellar pearlite was replaced by ausferrite when copper (Cu)¹⁰⁾ and nickel (Ni)¹¹⁾ were added to sintered Fe–Mo–Si–C alloys.

Due to multiple effects of Si addition as given above, this work was designed to explore the effects of alloying molybdenum (Mo), Si, and C elements on phase transformations under slow cooling after sintering of sintered Fe–Mo–Si–C alloys. It was expected that experimental results would lead to deep understanding on the combined effects of alloying Mo, Si and C elements on microstructural development and mechanical properties of sintered Fe–Mo–Si–C alloys.

2. Experimental Procedure

2.1. Material Processing

Three pre-alloyed powders were used to design samples for studying the combined effects of Mo, Si, and C by adding silicon carbide powder (Sigma-Aldrich Chemistry of the USA) in varying weight percentages of 1.0, 2.0, 3.0, and 4.0. The pre-alloyed powders consist of (1) pre-alloyed Fe-0.5Mo-0.15Mn powder (Rio Tinto Metal Powders of Canada), (2) pre-alloyed Fe-0.85Mo powder (Hoganas of Sweden), and (3) pre-alloyed Fe-1.50Mo powder (Hoganas of Sweden). By combining base powders with silicon carbide, we can prepare three sets of experimentally sintered alloys shown in Table 1. After weighing each powder, they were mixed and blended with a lubricant (1.0 wt.% zinc stearate) in a double-cone tumbling mixer for 3.6 ks. The tensile test bars were made by compacting these powder mixtures that met MPIF Standard 10. The density of the samples was determined using Archimedes’ method and the green density was maintained at 6.50 ± 0.05 Mg/m³. Sintering was performed at 1523 K for 2.7 ks in a vacuum furnace (Schmetz, Germany). The heating rate used during sintering was 0.16 K/s and was held at 873 K for 3.6 ks to remove the lubricant completely. The tensile test bars were subjected to slow cooling at a rate of 0.1 K/s after sintering. The selected cooling rate of 0.1 K/s is equivalent to the furnace cooling rate. The reasons for selecting this cooling rate include (i) the sufficient carbon diffusion during the cooling process can ensure the roles of other alloying elements on phase transformations involved and (ii) this cooling rate requires no cooling inert gas, leading to economic reason for potential production.

Table 1. Nominal composition in mixed raw powder and sintered density.

Alloy ID	Mixed powder (wt.%)		Nominal composition (wt.%)					Sintered density (g/cm³)
Alloy ID	Prealloyed Fe	SiC	C	Si	Mo	Mn	Fe	Sintered density (g/cm³)
05Mo00SiC	100	0	0.0	0.0	0.50	0.15	Bal.	7.21
05Mo10SiC	99	1	0.3	0.7	0.50	0.15	Bal.	7.13
05Mo20SiC	98	2	0.6	1.4	0.49	0.15	Bal.	7.08
05Mo30SiC	97	3	0.9	2.1	0.49	0.15	Bal.	7.16
05Mo40SiC	96	4	1.2	2.8	0.48	0.14	Bal.	7.63
08Mo00SiC	100	0	0.0	0.0	0.85	0.0	Bal.	7.45
08Mo10SiC	99	1	0.3	0.7	0.84	0.0	Bal.	7.38
08Mo20SiC	98	2	0.6	1.4	0.83	0.0	Bal.	7.35
08Mo30SiC	97	3	0.9	2.1	0.82	0.0	Bal.	7.24
08Mo40SiC	96	4	1.2	2.8	0.82	0.0	Bal.	7.63
15Mo00SiC	100	0	0.0	0.0	1.50	0.0	Bal.	7.43
15Mo10SiC	99	1	0.3	0.7	1.49	0.0	Bal.	7.41
15Mo20SiC	98	2	0.6	1.4	1.47	0.0	Bal.	7.35
15Mo30SiC	97	3	0.9	2.1	1.46	0.0	Bal.	7.28
15Mo40SiC	96	4	1.2	2.8	1.44	0.0	Bal.	7.62

2.2. Materials Characterization

Samples for microstructure examination of size about 6 × 6 mm were randomly extracted from the sintered samples. The sample cross-section was mounted in resin using a Struer LaboPress-3 hot mounting machine and cured at a temperature of 180°C for 5 minutes. To reveal the constituents and structure of metals, ASTM E3-11 guidelines for metallographic specimen preparation were followed. The grinding process started with 180 grits of SiC paper followed by 400, 600, 800, and 1200 grits, respectively. Each step must remove all scratches from the previous step. Polishing was performed using diamond abrasives of 6, 3, and 1 μm until a smooth and scratch-free surface was achieved. After that, the smooth surface of the samples was immersed in 2% nital etchant for 20 seconds and then cleaned with water and dried. The microstructure observation was done by using optical microscopy (Olympus STM7 microscopes) and scanning electron microscopes (FE-SEM, Hitachi SU5000).

To ascertain the uniform distribution of elements in sintered samples, electron probe micro-analyzer (EPMA) was employed to analyze chemical distribution in a sintered sample microstructure consisting of ferrite plates, carbide particles and martensite-austenite (MA) constituents. The EPMA machine employed was Shimadzu EPMA8050G.

To prepare the sample for X-ray diffraction (XRD) analysis, we first cut and smoothed the surface using wet SiC abrasive paper, starting with 80-grit and finishing with 400-grit. Dissolving the ferrite matrix in a reagent that would not attack the carbide phases was conducted. This process aimed to improve the clarity of the carbide XRD peaks. The step started with etching (40 seconds) with a 40% nitric acid solution in water and then cleaning (5 seconds) with a 50% hydrochloric acid solution in water, both at room temperature.¹²⁾ XRD analysis was performed using a Rigaku TTRAX III instrument equipped with a copper source (wavelength of 1.5406 Å). The specimens were scanned in the angular two-theta range from 30° to 100° with a scan speed of 0.05°/s and a step angle of 0.02°. Phase identification from XRD peaks was conducted by comparing them with known standards in the Joint Committee on Powder Diffraction Standards (JCPDS) card catalog.

2.3. Mechanical Properties

In this study, mechanical properties, including macrohardness and tensile properties, were studied as a function of microstructures and chemical composition. The 574 Series Wilson Rockwell Scale B Hardness Tester was used for macrohardness measurement. It was conducted on 15 randomly located points of the surface, and then the average values were reported. The tensile properties of the samples were tested at room temperature using an Instron 8801 universal testing machine following standard test procedures (ASTM E8) with a strain rate of 0.00174 s⁻¹ at room temperature. The ultimate tensile strength, yield strength, and elongation were obtained from the stress-strain curve.

3. Results

3.1. Sintered Alloys without SiC Addition

Microstructures of sintered 05Mo00SiC (Figs. 1(a) and 1(b)), 08Mo00SiC (Figs. 1(c) and 1(d)), and 15Mo00SiC (Figs. 1(e) and 1(f)) alloys showed large polygonal ferrite (PF) grains and pores at grain corners. With PF grains, there were some fine pits due to removal of unknown inclusions by etching action. The average ferrite grain size in these sintered alloys (Figs. 1(b), 1(d) and 1(f)) was about 50 μm. The grain size measurements were obtained using a linear interception method. Random photographs at 50X magnification were chosen. Random lines of known length were drawn, and the number of intercepts was then counted. XRD patterns of all sintered alloys without SiC addition showed only peaks corresponding to body-centered cubic (bcc) crystal structure of α-ferrite (Fig. 2).

Fig. 1. Microstructures of sintered 05Mo00SiC (a and b), 08Mo00SiC (c and d), and 15Mo00SiC (e and f) alloys. (Online version in color.)

Fig. 2. XRD patterns of sintered 05Mo00SiC, 08Mo00SiC and 15Mo00SiC alloys.

The transformation from austenite to ferrite in all sintered alloys without SiC addition is attributed to reconstruction phase transformation. The smooth grain boundaries obtained due to slow cooling rate of 0.1 K/s and equiaxed grain shape indicate that the phase transformation from austenite to ferrite starts with allotriomorph ferrite nucleation at prior austenite grain boundaries and grows into remaining austenite matrix. The average grain size of ferrite grains of about 50 μm is half of the D50 of Fe-0.85Mo powder (111.86 μm). This means that most grains are formed by impingement of two different allotriomorph ferrite crystals.

3.2. Sintered Alloys with 1.0 wt.% SiC Addition

Microstructures of sintered 05Mo10SiC (Figs. 3(a) and 3(b)), 08Mo10SiC (Figs. 3(c) and 3(d)), and 15Mo10SiC (Figs. 3(e) and 3(f)) alloys showed common microstructural feature consisting of PF grains, eutectoid transformation zones (ETZs) and pores at grain corners. Within each ETZ, there were coarse ferrite plates with thick or thin bands of ferrite + carbide (FC) mixture on coarse ferrite plate boundaries. Within a thin band of FC mixture, designated as FC1 mixture, ferrite and carbide showed degenerative growth mode with a row of needle carbide particles precipitating on ferrite plate boundaries. Within a thick band of FC mixture, designated as FC2 mixture, ferrite and carbide showed cooperative growth mode. The FC mixture in a thick band could be identified as pearlite.

Fig. 3. Microstructures of sintered 05Mo10SiC (a and b), 08Mo10SiC (c and d), and 15Mo10SiC (e and f) alloys.

According to ETZ components, the phase transformations in ETZ occur in two stages, such as the formation of coarse ferrite plates and the decomposition of austenite bands between coarse ferrite plates. Although the two-stage formation of ETZ has a similar manner to that of the two-stage formation of upper bainite given elsewhere.¹³⁾ According to the external morphologies of microstructurally defined bainite,¹⁴⁾ the upper bainite with ferrite plates of higher volume fraction preceding the formation of carbide of lower volume fraction may be a suitable terminology for the ETZ features, such as coarse ferrite plates plus FC1 mixtures and coarse ferrite plates plus FC2 mixtures.

XRD patterns of sintered 05Mo10SiC and 08Mo10SiC alloys showed strong peaks of α-ferrite and weak peaks of face-centered cubic (fcc) crystal structure of M₂₃C₆ carbide and orthorhombic crystal structure of M₃C carbide while that of sintered 15Mo10SiC alloy showed strong peaks of α-ferrite and weak peaks of γ-austenite and M₃C carbide (Fig. 4). Due to the presence of γ-austenite XRD peaks, some elongated bands on ferrite plate boundaries in sintered 15Mo10SiC could be retained austenite bands (Fig. 3(f)).

Fig. 4. XRD patterns of sintered 05Mo10SiC, 08Mo10SiC and 15Mo10SiC alloys.

With fixed 1.0 wt.% SiC addition, the size of PF grains decreased with increasing Mo content (Figs. 3(a), 3(c) and 3(e)). The distribution of ETZs was more highly uniform when Mo contents were increased (Figs. 3(b), 3(d) and 3(f)). The increase of Mo content also led to the formation of retained austenite bands on ferrite plate boundaries.

The decrease of large equiaxed PF grains is attributed to the increase of transformation undercooling or the decrease of transformation temperature. The equiaxed PF grains, formed at the highest temperatures and slowest cooling rates in low-carbon steels, are nucleated as grain-boundary allotriomorphs and grow into equiaxed grains.¹⁵⁾ As the transformation temperature is lowered below Widmanstätten ferrite start (W_S) temperature, Widmanstätten ferrite (WF) plates form instead of PF grains at temperatures around the nose of upper C curve of the time-temperature-transformation (TTT) diagram.¹⁶⁾

It is commonly known that Mo addition favors the generation of bainite at relatively low cooling rates and promotes a finer microstructure.¹⁷⁾ The promotion of WF plates in sintered 08Mo10SiC and 15Mo10SiC alloys is attributed to the Mo role in delaying phase transformation from face-centered cubic (fcc) to body-centered cubic (bcc) and promoting nonpolygonal ferrite formation with high dislocation density.¹⁸⁾ The role of Mo on acicular ferrite formation was found to comprise two aspects, such as affecting transformation kinetics and modifying the morphology of the acicular ferrite.¹⁹⁾

The presence of retained austenite bands between WF plates indicates the combined effect of Mo, Si and C on austenite stabilization. In the previous work, sintered Fe-1.50Mo-0.30C alloy, produced without Si addition and under sintering furnace cooling, showed only PF grains and ETZs without retained austenite.²⁰⁾ When Si is introduced to sintered 05Mo10SiC alloy (with 0.50 wt.% Mo), the combined effect of alloying Mo, Si and C elements on austenite stabilization is insufficient to produce retained austenite bands on WF plate boundaries. When the Mo content is 1.50 wt.% in sintered 15Mo10SiC alloy, the combined effect of alloying Mo, Si and C elements on austenite stabilization is effective.

3.3. Sintered Alloys with 2.0 wt.% SiC Addition

The OM microstructures of sintered 05Mo20SiC (Fig. 5(a)), 08Mo20SiC (Fig. 5(c)), and 15Mo20SiC (Fig. 3(e)) alloys showed only ETZs consisting of Widmanstätten ferrite (WF) plates with ferrite + particle (FP) mixtures on WF plate boundaries. With close observation by SEM, it was revealed that the microstructural components in ETZs were different among these sintered alloys. The ETZ of sintered 05Mo20SiC alloy comprised WF plates with thick and thin FC mixture bands on WF plate boundaries (Fig. 5(b)). In sintered 08Mo20SiC alloy (Fig. 5(d)), a new feature of FP mixture, consisting of alternating ferrite and austenite plates and assigned as degenerate upper bainite (DUB) as given elsewhere,²¹⁾ was observed on WF plate boundaries. The DUB units in sintered 08Mo20SiC alloy (Fig. 5(d)) were confirmed by moderate XRD peaks of γ-austenite (Fig. 6).

Fig. 5. Microstructures of sintered 05Mo20SiC (a and b), 08Mo20SiC (c and d) and 15Mo20SiC (e and f) alloys. (Online version in color.)

Fig. 6. XRD patterns of sintered 05Mo20SiC, 08Mo20SiC and 15Mo20SiC alloys.

The microstructural feature in sintered 15Mo20SiC alloy (Fig. 5(f)) was different from that of sintered 08Mo20SiC alloy (Fig. 5(d)). The WF plates disappeared from the former sintered alloy. The thicknesses of ferrite and austenite plates in sintered 15Mo20SiC alloy (Fig. 5(f)) were roughly equal. Since this microstructural feature was like that of upper ausferrite reported in an austempered ductile iron (ADI),²²⁾ it was identified as upper ausferrite (UA).

3.4. Sintered Alloys with 3.0 wt.% SiC Addition

The microstructures of sintered 05Mo30SiC (Figs. 7(a) and 7(b)), 08Mo30SiC (Figs. 7(c) and 7(d)), and 15Mo30SiC (Figs. 7(e) and 7(f)) alloys drastically changed from steel to cast iron microstructural feature. Ductile iron (DI)-like microstructures were observed in these sintered alloys. In sintered 05Mo30SiC alloy (Figs. 7(a) and 7(b)), the common microstructural feature comprised black particles embedded in a matrix of ferrite halo and pearlite nodules. The black particle was formed by solidification of a melt generated around the SiC particle site due to eutectic melting.¹¹⁾ The solid graphite formed either on a pore surface or on a solid SiC residue surface. When solidification completed, graphite and austenite coexisted and followed by stable eutectoid transformation of austenite to ferrite halo and metastable eutectoid transformation of austenite to pearlite.⁸⁾ This type of microstructural feature resembled that of a ferrite-pearlite DI. In sintered 08Mo30SiC (Figs. 7(c) and 7(d)) and 15Mo30SiC (Figs. 7(e) and 7(f)) alloys, the common microstructural feature comprised black particles embedded in a matrix of ferrite halo, pearlite and ausferrite. This type of microstructural feature has been rarely reported in DIs.

Fig. 7. Microstructures of sintered 05Mo30SiC (a and b), 08Mo30SiC (c and d) and 15Mo30SiC (e and f) alloys. (Online version in color.)

The carbides associated with pearlite nodules in sintered 05Mo30SiC, 08Mo30SiC, and 15Mo30SiC alloys were different as indicated by XRD patterns given in Fig. 8. Mixed M₃C and M₂₃C₆ carbides were observed in the pearlite nodules of sintered 05Mo30SiC alloy while M₂₃C₆ became dominant in the pearlite nodules of sintered 08Mo30SiC and 15Mo30SiC alloys.

Fig. 8. XRD patterns of sintered 05Mo30SiC, 08Mo30SiC and 15Mo30SiC alloys.

3.5. Sintered Alloys with 4.0 wt.% SiC Addition

The microstructures of sintered 05Mo40SiC (Figs. 9(a) and 9(b)), 08Mo40SiC (Figs. 9(c) and 9(d)), and 15Mo40SiC (Figs. 9(e) and 9(f)) alloys were like an iron microstructure. DI-like microstructures were observed in these sintered alloys. In sintered 05Mo40SiC (Figs. 9(a) and 9(b)) and 08Mo40SiC (Figs. 9(c) and 9(d)) alloys, the common microstructural feature comprised black particles embedded in the matrix of ferrite halos and pearlite nodules. The carbide associated with pearlite in sintered 05Mo40SiC and 08Mo40SiC alloys was M₃C type, as given in corresponding XRD patterns (Fig. 10).

Fig. 9. Microstructures of sintered 05Mo40SiC (a and b), 08Mo40SiC (c and d) and 15Mo40SiC (e and f) alloys. (Online version in color.)

Fig. 10. XRD patterns of sintered 05Mo40SiC, 08Mo40SiC and 15Mo40SiC alloys.

In sintered 15Mo40SiC (Figs. 9(e) and 9(f)) alloy, the common microstructural feature comprised black particles embedded in the matrix of ferrite halo, pearlite, and ausferrite. The carbides associated with pearlite in sintered 15Mo40SiC alloy included M₃C and M₂₃C₆ types, as given in corresponding XRD patterns (Fig. 10). The moderate XRD peaks of γ-austenite in sintered 15Mo40SiC alloy (Fig. 10) confirmed the formation of ausferrite in this sintered alloy. This type of microstructural feature, regarding the coexistences of pearlite and ausferrite and that of M₃C and M₂₃C₆ carbides, was reported in a few DIs. The coexistence of pearlite and ausferrite was reported in some as-cast gray irons.²³⁾

In sintered Fe–Mo–Si–C alloys, produced from different Fe–Mo powders mixed with 4.0 wt.% SiC, showed microstructural change in matrices with Mo content, i.e., the matrix microstructures changed from ferrite halo plus pearlite to ferrite halo plus pearlite and ausferrite with increasing Mo content.⁹⁾ The combination of high Mo, Si, and C contents was suspected to contribute to ausferrite formation. The characterization results given in Fig. 10 confirmed the previous study results.⁹⁾

3.6. Mechanical Properties of Experimental Sintered Alloys

This experimental work was conducted to find the relative effect of SiC content on mechanical properties of experimental sintered alloys. It was found that the increase of added SiC content led to increases in tensile strength and hardness (Fig. 11). The usual strength and ductility trade-off was observed when SiC contents of up to 2.0 wt.% were added. The synergy of strength and ductility was observed in the cases of added SiC contents were 3.0 and 4.0 wt.%.

Fig. 11. Mechanical properties of experimental sintered alloys. The symbols 05Mo, 08Mo and 15Mo represent pre-alloyed Fe-0.5Mo-0.15Mn, Fe-0.85Mo, and Fe-1.50Mo powders, respectively.

Taking microstructures of sintered alloys with steel microstructures into consideration, it is found that strength and hardness increase whereas the ductility decreases when the fractions of FC mixtures, DUB, and ausferrite are increased. When DI-like microstructures are developed in sintered alloys, the synergy of strength and ductility can be obtained. This synergy may be attributed to the pearlite structure, the dominant microstructural component of sintered DI-like alloys. The yield stress, the ultimate tensile strength and the ductility exhibit a rising trend with smaller pearlite interlamellar spacing and a lower pearlitic colony size.²⁴⁾ The mechanisms for strength and ductility synergy by lamellar structure are given elsewhere.²⁵⁾

4. Discussion

4.1. Combined Effects of Mo, Si and C on Austenite Stabilization

Before discussing on the combined effects of Mo, Si and C on austenite stabilization, it should be noted here that distributions of alloying elements are not always uniform in sintered steels. To ascertain the uniformity of the elemental distribution of Mo, Si and C in a sintered alloy, EPMA was conducted on microstructure of sintered 08Mo20SiC alloy, consisting of ferrite plates, carbide particles and martensite-austenite (MA) constituents (Fig. 12).

Fig. 12. Chemical distribution in microstructure of sintered 08Mo20SiC alloy, consisting of ferrite plates, carbide particles, and martensite-austenite constituents. (Online version in color.)

The EPMA results suggest that the distributions of Fe, Mo, Si, Mn are uniform except that of C has high concentrations in carbide particles and MA constituents. The C distribution in carbide particles shows heterogeneity compared to the homogeneous C distruibution in MA constituents (Fig. 12(f)). The EPMA results also suggest that alloying C element has strong influence on austenite stabilization as its high concentration remains in MA constituents at room temperature.

The presence of austenite phase in slowly cooled sintered alloys, such as 15Mo10SiC (Figs. 3(f) and 4), 08Mo20SiC (Figs. 5(d) and 6) and 15Mo20SiC (Figs. 5(f) and 6) alloys, is a very interesting issue for discussion here. When the compositions of these sintered alloys are considered, only alloying C element acts as austenite stabilizer. The nominal C content in sintered 15Mo10SiC alloy is 0.3 wt.%, which is questionable for austenite stabilization. However, the C partitioning from ferrite plates to enrich adjacent untransformed austenite bands during austenite to ferrite transformation can occur and austenite stabilization by C enrichment becomes possible. According to previous work,²⁶⁾ the microstructure of sintered Si-free Fe-1.5Mo-0.3C alloy had only polygonal grains and ETZs without retained austenite. Thus, the presence of austenite in sintered 15Mo10SiC alloy can be attributed to alloying Si element. However, Si is not a direct austenite stabilizer. Its role in austenite stabilization is not straightforward. The judicious Si additions can hinder carbide precipitation in austenite³⁾ due to thermodynamic reasons.^4,6) When carbide precipitation is retarded, austenite can be stabilized by solute C atoms.⁶⁾ By the principle given above, Si is claimed to enhance thermal stability of austenite in some previous works.²⁷⁾

In sintered alloys with 2.0 wt.% SiC addition, it is revealed that retained austenite is not observed in sintered low Mo-containing 05Mo20SiC alloy but clearly observed in sintered high Mo-containing 08Mo20SiC and 15Mo20SiC alloys (Fig. 6). The austenite amounts, indicated by the XRD peak intensities of γ-austenite in sintered 08Mo20SiC and 15Mo20SiC alloys (Fig. 6), also increase with increasing Mo content. This seems to indicate that the alloying Mo element also has a role in austenite promotion. However, the same as alloying Si element, Mo is not a direct austenite stabilizer. Mo involvement in austenite promotion is not considered as its direct function. According to previous studies, the Mo effect on retained austenite amount was controversial. In an austempered ductile iron (ADI), the retained austenite content increases with the content of copper, decreases with the content of molybdenum, and reaches the maximum with a certain content of silicon.²⁸⁾ In contrast, it was demonstrated that the Mo addition in low carbon steels led to enhancement of retained austenite formation under the carburized and hardened conditions.²⁹⁾ However, the exact action of Mo was not given in this work. An increase in the amount of alloying element, i.e., Mo, is accompanied by an increase in the tendency to form retained austenite.³⁰⁾

Mo can have indirect effects on austenite stabilization. Chen³¹⁾ reported that two metallurgical functionalities of Mo are: firstly, adding molybdenum partially inhibits the development of pearlite but significantly impedes the formation of ferrite. Secondly, carbon-enriched austenite is difficult to transform into bainite if molybdenum is added. However, in heavy-section castings, the addition of less than 1.0% molybdenum is not enough to prevent the production of pearlite. This is because the phase field of pearlite in the continuous cooling transformation (CCT) diagram is moved toward longer times. The pearlite suppression in steels is evidenced in several research works.^18,32) The bypass of pearlite transformation by Mo leads to the preservation of austenite down to lower temperatures, at which low temperature phase transformations can occur. The Mo-containing steel with composition of Fe-0.22C-1.57Si-2.00Mn-0.026Nb-0.14Mo (wt.%) produced under continuous cooling at 0.2 K/s had microstructure consisting of bainite-martensite instead of ferrite-pearlite microstructure in Mo-free steel with composition of Fe-0.22C-1.55Si-2.01Mn-0.025Nb (wt.%) produced under the same cooling rate.³³⁾

4.2. Pearlite Transformation due to High C Concentration

In the sintered 05Mo30SiC alloy (Figs. 7(a) and 7(b)), the formation of ferrite halos and lamellar pearlite nodules around black particles makes its microstructure resemble that of a ferrite-pearlite ductile iron. In sintered 08Mo30SiC (Figs. 7(c) and 7(d)) and 15Mo30SiC (Figs. 7(e) and 7(f)) alloys, the ausferrite coexist with ferrite halos, pearlite nodules, and black particles. Since Mo is commonly known to have the effect on pearlite transformation retardation the pearlite transformation in these sintered Mo-containing steels is worth discussing here.

The carbon content in carbon steels controls the discontinuous nature, thickness, and volume fraction of cementite in pearlite.³⁴⁾ According to Yasuda and Nakada,³⁵⁾ the cementite morphology in pearlite changed from a lamellar to a spherical shape because of a reduction in the carbon concentration in the austenite. At a temperature of 773 K, the critical carbon concentration for this transition in cementite morphology is approximately 0.42 wt.%. Thus, lamellar pearlite prefers to form from austenite with a high C concentration.

In sintered Mo-containing high C steels, the formation of lamellar pearlite is difficult. Sintered Fe-0.50Mo-0.15Mn-0.90C alloy had a microstructure mainly consisting of ferrite + needle M₃C carbide mixtures without conventional lamellar pearlite feature.³⁶⁾ Surprisingly, the 2.0 wt.% SiC addition leads to the formation of cooperative ferrite + carbide mixtures (ferrite + M₂₃C₆ and ferrite + M₃C) with features close to lamellar pearlite in thick austenite bands between WF plates in sintered 05Mo20SiC (Figs. 5(a), 5(b) and 6). With the 3.0 wt.% SiC addition, M₂₃C₆ and M₃C lamellae form as lamellar pearlite components in the sintered 05Mo30SiC alloy (Figs. 7(a), 7(b) and 8) while dominant M₂₃C₆ lamellae are found in lamellar pearlite structures of sintered 08Mo30SiC (Figs. 7(c), 7(d) and 8) and 15Mo30SiC alloys (Figs. 7(e), 7(f) and 8). It is seeming that the Si role in carbide precipitation retardation is weakened by high C concentration. As given by Kozeschnik et al.,⁶⁾ if the carbon in the parent phase is highly supersaturated, Si does not influence the precipitation of cementite. The reason for this is that even in the case of para-equilibrium precipitation, the reaction has a strong driving force.

4.3. Formation of Ausferrite in Slowly Cooled Sintered Fe–Mo–Si–C Alloys

The most striking phenomenon of microstructural developments is the formation of upper ausferrite feature in some slowly cooled sintered Fe–Mo–Si–C alloys. Ausferrite is the common microstructural feature obtained from the austempering process, which is conducted after quenching from austenitization temperatures at low isothermal temperatures, at which bainitic ferrite plate formation leads to C enrichment of adjacent austenite plates.³⁷⁾

In some experimental sintered alloys given above, the ausferrite formation occurs via slow and continuous cooling. The ausferrite formation in slowly cooled sintered alloy must be associated with austenite stability mentioned above (Topic 4.1) and occurs at the right phase transformation temperature range. The dominant upper ausferrite feature in slowly cooled sintered 15Mo20SiC alloy (Figs. 5(e) and 5(f)) indicates the right composition and temperature range window for ausferrite formation. The coexistence of upper ausferrite with ferrite halos and pearlite nodules in slowly 08Mo30SiC alloy (Figs. 7(c) and 7(d)), 15Mo30SiC alloy (Figs. 7(e) and 7(f)), and 15Mo40SiC alloy (Figs. 9(e) and 9(f)) indicates the cascading phase transformations of ferrite, pearlite and ausferrite in these sintered alloys.

5. Conclusions

The microstructural evolution and mechanical properties of sintered alloys produced by varied SiC additions to pre-alloyed Fe–Mo powders with different Mo contents were revealed in this study. With 1.0 wt.% SiC addition, slowly cooled sintered alloys showed microstructures consisting of PF grains and ETZs in the form of upper bainite. With 2.0 wt.% silicon carbide addition, sintered alloys with Mo contents of ≥ 0.85 wt.% had microstructures consisting of DUB or ausferrite. With ≥ 3.0 wt.% silicon carbide addition, DI-like microstructures were developed in sintered alloys. The change of microstructures in experimental sintered alloys was attributed to the combined effect of alloying molybdenum, silicon, and carbon elements. Tensile strength and hardness increased with increasing added SiC content while ductility varied with microstructural components.

Statement for Conflict of Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgments

The authors would like to thank all support for this work, which is financially supported via the project ‘Design and manufacturing of replacement parts for railway applications (P1951261)’ under NSTDA, Pathum Thani, Thailand. Technical supports are obtained from National Metal and Materials Technology Center (MTEC), Pathum Thani, Thailand.

References

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